Boosting the interfacial superionic conduction of halide solid electrolytes for all-solid-state batteries

Designing highly conductive and (electro)chemical stable inorganic solid electrolytes using cost-effective materials is crucial for developing all-solid-state batteries. Here, we report halide nanocomposite solid electrolytes (HNSEs) ZrO2(-ACl)-A2ZrCl6 (A = Li or Na) that demonstrate improved ionic conductivities at 30 °C, from 0.40 to 1.3 mS cm−1 and from 0.011 to 0.11 mS cm−1 for Li+ and Na+, respectively, compared to A2ZrCl6, and improved compatibility with sulfide solid electrolytes. The mechanochemical method employing Li2O for the HNSEs synthesis enables the formation of nanostructured networks that promote interfacial superionic conduction. Via density functional theory calculations combined with synchrotron X-ray and 6Li nuclear magnetic resonance measurements and analyses, we demonstrate that interfacial oxygen-substituted compounds are responsible for the boosted interfacial conduction mechanism. Compared to state-of-the-art Li2ZrCl6, the fluorinated ZrO2−2Li2ZrCl5F HNSE shows improved high-voltage stability and interfacial compatibility with Li6PS5Cl and layered lithium transition metal oxide-based positive electrodes without detrimentally affecting Li+ conductivity. We also report the assembly and testing of a Li-In||LiNi0.88Co0.11Mn0.01O2 all-solid-state lab-scale cell operating at 30 °C and 70 MPa and capable of delivering a specific discharge of 115 mAh g−1 after almost 2000 cycles at 400 mA g−1.

11. In Fig. 4, the authors claimed that "Li2ZrF6 and Li3Zr4F19 were thermodynamically stable passivating interphases that were decomposed from Li2ZrCl5F.  (1) I am wondering whether the LiF can be decomposed from Li2ZrCl5F. Why or why not?
13. "cost-effective and abundant elements, such as Zr and Al" (Line 415) I don't think "Zr" is an "abundant" element although authors compared ZrCl4 with other rare metal chlorides. There might be more appropriate words.
14. Please revise the wrong spell of 'Thoertical calculations'(Line 588) Reviewer #2 (Remarks to the Author): This paper reports interfacial superionic conduction in ZrO2-Li2ZrCl6 system. This paper is interesting, but it needs to address the following questions: (1) It is unclear that why nZrO2-Li2ZrCl6 has lower conductivity, since Li2ZrCl6 should also have a percolation network.
(4) What is the interacton between F and O in ZrO2-LZCF system? Reviewer #3 (Remarks to the Author): The manuscript by Kwak targeted the preparation of novel halide nanocomposite electrolyte compositions with general chemical formula ZrO2-AX-A2ZrX6, with A = Li or Na, X = Cl, F, which are claimed to have good interfacial ionic conductivity as well as good stability against sulfide solid electrolytes. Furthermore, the author claim this new class of electrolytes appears stable against LiCoO2 and high-voltage NMC electrodes under fast charging, relatively high-cycling temperatures (60 Celsius), and long-term cycle life. The mechanism of interfacial ion transport has been assessed using density functional theory. The validity of these mechanisms is further strengthened by magic-angle spinning solid-state NMR and specialized synchrotron experiments. Nicely, DFT could verify that oxygen "doping" on the Cl sites, increased the Zr-anion bond length, in agreement with the EXAFS analysis.
The manuscript is interesting and should be considered by Nature Communication after my comments are addressed.
1. The vendor and purity of the materials utilized in the synthesis of the halide nanocomposite should be reported.
5. High-resolution transmission electron microscopy is known to damage Li2YCl6. I must admit that the micrographs in Figure S3 are not clear, and the diffraction not neat. I wonder whether Li2ZrCl6 derivatives can be equally damaged by the electron beam. The authors should clarify this.
6. Furthermore, the claim "Interestingly, the HRTEM images of the ZrO2-2Li2ZrCl6 HNSE (Figure 1f,g) showed that ZrO2 formed a percolating network nanostructure with thickness of only a few nanometres." Does not seem supported by the data. If they don't have data backing this statement, this should be removed. 7. The statement "To the best of our knowledge, the Li+ HNSE was the first inorganic superionic conductor that exploited the interfacial effect to promote conduction with ionic conductivity reaching 1 mS cm-1." appears a bit over the top. To me this is just the result of ZrO2 modifying the texture of the microstructure, i.e. grain boundary, and hence improving the ion transport. These types of claims should be toned down.
8. Furthermore, it seems that many compounds reported in Tables S3 and S4 have a sizeable amount of both ZrO2 and LiCl (or NaCl for the Na analogues) and are far from being phase pure. Now, this study clearly glosses over the importance of these impurities on the texture of grain boundaries, and thus the conductivity properties. A SEM analysis of the most important component is required. There may be a significant statistic of particle size which may also impact these observations. 9. Figure 2d, it's unclear why there is just a datapoint for nZrO2-2Li2ZrCl6. Why couldn't the author measure the ion conductivity at lower or higher temperatures? 10. In the supplementary material, I don't see any detail about the impedance analysis, starting from the equivalent circuits that have been used to extract the claimed ionic conductivities. This level of analysis is paramount. In particular, it's not clear whether they're able to deconvolute the grain vs bulk contribution to the total conductivity. Nice semi-circles are visible for the Na-based compounds. Have they tried? 11. The authors claimed that "Li+ diffusion was ~11 times faster for LZCO than LZC at 300 K". This is a very big claim considering that AIMD at a temperature as low as 300 cannot be considered converged. Perhaps is more valuable to make this comparison at higher temperatures.
Minor comments: 1. In the introduction, it should be "oxidative limits" not "oxidation limits".
2. In the introduction, It's true that oxide solid electrolytes are brittle, but this is not the main challenge. The main challenge in the utilization of oxide-based solid electrolytes is the cost-intensive process of sintering these materials, which is required at either very high temperatures or highly specialized systems such as spark plasma sintering.

3
(2) "The XRD signals of ZrO2 were not observed, suggesting nanosized grains or poor crystallinity" (Line 121). I don't think the no observation of the XRD peak of ZrO2 is due to nanosized grains, according to the Scherrer equation.
Response to comment 2: We are thankful to the reviewer for the insightful comments.
(1) According to the comment, we have conducted a complementary analysis using DFT calculations, synchrotron XRD and PDF measurements. First, through DFT calculations, the chemical reaction energies are calculated using the energies of reactants and products (Supplementary Table 1). The reaction to generate ZrO2 and LiCl from ZrCl4 and Li2O has a strong driving force (Equation S1, ΔE = -4.736 eV). Furthermore, there is another spontaneous reaction from reactants (Li2O, ZrCl4, LiCl) to products (Li2ZrCl6, ZrO2) when the molar ratios are stoichiometrically matched (Equation S2, ΔE = -5.000 eV). Therefore, it is energetically not favorable for Li2O to remain as an impurity after the reaction. Moreover, ternay-compound phase diagrams of ZrO2-ZrCl4-LiCl and Li-Zr-Cl are plotted in Supplementary Figure 2. No stable compounds exist except for Li2ZrCl6, indicating that the formation of impurity is energetically less favorable than one of Li2ZrCl6 and ZrO2 although local off-stoichiometry may occur during the reaction. Furthermore, to verify the composition of ZrO2-2Li2ZrCl6, synchrotron XRD and PDF measurements were carried out for a precursor mixture of Li2O and ZrCl4 (2:3 molar ratio) with varying ball-milling time (Supplementary Figures 3, and 4). The XRD patterns with increasing ball-milling time are classified into three regions ( Supplementary Fig. 3). In region I corresponding to 2, 6, and 8 h, the peaks for LiCl emerge with slightly decreased peak intensities for ZrCl4 and Li2O. In region II (10, 11, 12, and 16 h), the Li2ZrCl6 peaks evolve at the expense of the lowered peak intensities for ZrCl4, Li2O, and LiCl. Finally, after 20 h (region III), the XRD patterns remain almost identical with the exception of a minor decrease in the peak for Li2O.

Supplementary
Notably, for the sample ball-milled for 30 h, only the Li2ZrCl6 peaks are present with no detectable impurity or precursor peaks. Consistently, the PDF analysis results also show that, as the ball-mill time increases, the signal intensity for Li2ZrCl6 increases at the expense of the lowered intensities of the Li2O and ZrCl4 signals (Supplementary Figure 4). Moreover, the PDF analysis results confirm that the amount of Li2ZrCl6 does not increase further after 20 h. In addition, the evolution and increase of a ZrO2 peak around 2 Å, corresponding to Zr-O bonding, is corroborated during the mechano-chemical milling.
In summary, the complementary analysis results unequivocally lead us to the following conclusions. The synthesis of ZrO2 and Li2ZrCl6 from the reaction between Li2O and ZrCl4 is energetically favorable. The mechanochemical synthesis of ZrO2-2Li2ZrCl6 HNSE without a negligible amount of precursors or impurities is confirmed when enough time (≥ ≈20 h) and energy are provided.
In the revised manuscript, Supplementary Figures 2, 3, and 4, and Supplementary Table 1 have been added. The relevant discussion has also been added in the Section of "Synthesis and characterization of HNSEs". Also, the expression about the composition has been revised as follows.
"We extensively characterized the ZrO2-2Li2ZrCl6 HNSE sample as it is a simple binary system and exhibited a much higher Li + conductivity of 1.1 mS cm -1 than Li2ZrCl6 (0.40 mS cm -1 ), despite the 7.86 vol.% of insulating ZrO2 (based on the chemical formula of ZrO2-2Li2ZrCl6)." 5 Supplementary Fig. 3 Characterization of the synthesis reaction mechanism for HNSE by Synchrotron XRD. Synchrotron XRD patterns for the precursor mixture of Li2O and ZrCl4 (2:3 molar ratio) with varying ball-milling time.
6 Supplementary Fig. 4 Characterization of the synthesis reaction mechanism for HNSE by PDF. Synchrotron PDF G(r) for the precursor mixture of Li2O and ZrCl4 (2:3 molar ratio) as a function of ball-milling time.
(2) According to the HRTEM images, ZrO2 in HNSEs exists as nanosized grains (~ 10 nm) with poor crystallinity (Figures 1f,g). Please read our response to the Reviewer 3's response 5 for the details. Besides, the PDF spectra reveal the ZrO2 peaks up to ≈10 Å (Supplementary Figure 4), indicating that amorphous ZrO2 is also present in HNSEs. R1 In conclusion, ZrO2 in HNSEs exists as both crystalline and amorphous phases, while the latter occupies a larger fraction. The absence of ZrO2 peaks in the XRD patterns (Figures 1b, and Supplementary Fig. 2) is thus understood.
In the revised manuscript, the term has been toned down as the following: "nanosized grain" "nanosized grain with poor crystallinity". Response to comment 3: We appreciate the reviewer pointing out the possibility of remaining impurities in the final product. To verify the composition of the final product, we performed extensive PDF fitting refinements for ZrO2-2Li2ZrCl6 (prepared by ball-milling for 20 h) using various combinations of structure models, including the LZCO interphase structure provided via DFT calculations, as indicated in the revised Supplementary Table 7. After such extensive preliminary refinements, the best-fit results are obtained using two different refinement ranges; low r range of 1.5−10 Å (region 1), high r range of 10−30 Å (region 2). In all cases, the medium and average structure of ZrO2-2Li2ZrCl6 in the high r range (region 2) can be well represented by using a single Li2ZrCl6 structure (Supplementary Figure 17 and Table 10). For the low r range 1.5 ~ 10 Å (region 1), where the interface regime becomes prevailing, thus we have tested various combinations of model structures, including the precursors (e.g., LiCl and ZrCl4), Li2.5ZrCl5.5O0.5 (LZCO) interphase provided via DFT calculations, and Li2O impurity, as tabulated in the supplementary Table 10. Remarkably, we achieved the excellent reliability factor and best-fit result when the significant amount of theoretically suggested LZCO interphase and minor Li2O impurity are included (i.e., composition 3 in Supplementary Table 10). Please note that the reliability factor Rw of 48.3% in the case of using only two ZrO2 and Li2ZrCl6 model structures in Fig. 1d is significantly lowered to 10.9% in the case of including the LZCO interphase. The calculated composition obtained by the PDF fit corresponds to 1.47Li2ZrCl2-0.36Li2.5ZrCl5.5O0.5(interphase)-1.01ZrO2-0.16Li2O (Supplementary Table 11). This result strongly supports the abundant LZCO interphase existing between ZrO2 and LZC nanodomains.
We also performed the PDF fit for 30 h ball-milled ZrO2-2Li2ZrCl6, which showed the best-fit result using the combinations of the same model structures as in the 20 h case. ( Figure R1). Notably, the calculated composition of the 30 h sample showed significantly decreased impurity Li2O amount (Table R1). Therefore, the Li2O impurity could be eliminated from the final products if the mechanochemical reaction time is sufficient enough.
We also note that the inclusion of precursor LiCl and ZrCl4 phases in the ZrO2-2Li2ZrCl6 results in unreliable fitting results and increasing Rw during the PDF refinement, as shown in Supplementary  Table 10. Therefore, we can unambiguously rule out the remaining precursors after sufficient ballmilling time for the ZrO2-2Li2ZrCl6.
In the revised manuscript, Supplementary Figure 17 and Supplementary Tables 10 and 11 have been added. The relevant discussion has also been added in the "Interfacial Superionic Conduction of HNSEs" section.
8 Supplementary Fig. 17 Experimental PDF with best-fit results for ZrO2-2Li2ZrCl6 in the 1.5−30 Å range. The PDF fitting was performed across different refinement ranges (low r range of 1.5−10 Å; high r range of 10−30 Å).

Supplementary Table 10
Reliable factor and mass fraction change obtained from PDF refinement in 1.5−10 Å with various compositions of precursors and products for ZrO2-2Li2ZrCl6.

Response to comment 4:
We are thankful to the reviewer for the insightful comment. As discussed in our response to Comment 2, theoretically, the reaction of Li2O with ZrCl4 to form LiCl and ZrO2 is spontaneous, which has also been verified experimentally by the XRD results shown in Figure R2. In the revised manuscript, a discussion about the reaction mechanism has been added.

Fig. R2 Two-step preparation of ZrO2-2Li2ZrCl6
HNSEs. XRD patterns for the product formed by the reaction of Li2O and ZrCl4 (top, step I) and the product formed by the subsequent reaction with ZrCl4 (bottom, step II). The formation of LiCl and Li2ZrCl6 by step I and II is verified, respectively. Fig. S14. Please specify the details of fluorination in the manuscript.

Comment 5. The authors claimed the significance of "fluorinated HNSE ZrO2-2Li2ZrCl5F". However, the specific approach of fluorination is not detailed in the part of "Synthesis and characterization of HNSEs" in the manuscript, and only a simple preparation protocol is exhibited in
Response to comment 5: We are thankful to the reviewer for the careful comment. According to the comment, the synthesis procedure of the fluorinated HNSE ZrO2-2Li2ZrCl5F has been added in the section of method (Preparation of materials) in the revised manuscript as the following.
" Response to comment 6: We are thankful to the reviewer for the comment. According to the comment, all-solid-state cells with higher mass loadings (22.4 and 9.9 (mg of cathode material) cm -2 for LCO and S-NCM88, respectively) were tested and the results are shown in Figure R3. The LCO electrode using ZrO2-LZCF with the higher mass loading (20.4 mg cm -2 ) exhibited an outstanding cycling performance at 60 o C, which is consistent with the results for the lower mass loading (7.7 mg cm -2 ) ( Figure R3a,b). However, the LCO electrode using ZrO2-LZC showed a poorer capacity retention for the higher mass loading, compared with that for the lower mass loading. In addition, the rate capability results at 30 o C for the S-NCM88 electrode using ZrO2-LZCF shows a significant degradation by increasing the mass loading from 3.7 to 9.9 mg cm -2 . Although the HNSE approach could counterbalance the degradation of ionic conductivity, the conductivity of ZrO2-LZCF, 0.49 mS cm -1 , thus needs to be enhanced further for fulfilling the requirement of practical applications. Furthermore, engineering for composite electrodes such as dry coating of halide SEs could work to enhance the rate capability by maximizing ionic contacts with cathode active materials (CAMs). R2 These perspectives are included in our future research.
In the revised manuscript, the thickness (and mass loading) of the separating SE layers and Li-In counter electrodes has been added.  Response to comment 8: We are thankful to the reviewer for the comment. As pointed out, the electrochemo-mechanical issues are critical for ASSBs. R3-7 For cathodes using halide SEs, side reactions between CAMs and SEs are marginal at mild conditions (e.g., cutoff voltage of 4.3 V and room temperature). R4,R8 Thus, the volume changes of CAMs during charge and discharge are the primary reason for the electrochemo-mechancial degradation in ASSBs. The disintegration of secondary particles of CAM and CAM/SE contact losses deteriorate Li + and e − transport pathways. R4-6 Especially, secondary particles of polycrystalline high-Ni NCM materials consisting of randomly oriented grains undergo severe disintegration in ASSB cells even at the initial cycles, leading to the poor initial Coulombic efficiency and fast capacity fading upon repeated cycling. R4-6 In our previous study, we demonstrated that, compared to polycrystalline CAMs, single-crystalline CAMs show better mechanical integrity and exhibited the excellent cycling stability after 100 cycles. R4 In this regard, to minimize the electrochemo-mechanical effects, the mechanochemically prepared HNSEs were applied for the single-crystalline CAMs of LiCoO2 and NCA88 without any protective coating layers. In addition, a high operating pressure of 70 MPa helps suppress the electrochemo-mechanical degradation in ASSB cells. For practical application, an assessment under low pressures of a few megapascals at maximum and corresponding engineering, such as developments of zero-strain CAMs, advanced binders with improved mechanical strength, and highly deformable SEs, is imperative. R3,9-14 However, those are out of the scope in this study.

Comment 9.
Is there a problem with gas emissions in the full battery? Why or why not?
Response to comment 9: We are thankful to the reviewer for the comment. NCM cathode materials are known to cause gas evolution when used in liquid electrolyte-based cells. R15,16 In recent literature, an observation of the gas evolution in ASSB cells made of NCM and sulfide SE was reported. R17,18 The structure of NCM is unstable at a deeply delithiated state due to the formation of significant amounts of highly reactive Ni 4+, which is prone to reduction to form the rock-salt-like phase on the surface. R5,16,19 The layered-to-rocksalt phase transformation process usually induces irreversible oxygen release from the surface of NCM. R19 Therefore, irreversible oxygen release can be prevented if the structural changes are suppressed.
In our previous result, microstructural evolution of NCM in ASSBs using sulfide or halide SEs was investigated. 40 In the NCM electrode employing sulfide SEs, after 100 cycles, a 20-30 nm thick surface layer that was distinct from the core region was observed. This layer was characterized as the NiO-like rocksalt structure. By contrast, NCA electrodes employing halide SEs did not show any difference in the crystal structure between the core and surface regions after 100 cycles. It is thus expected that there will be negligible gas evolution in ASSBs cells with NCA electrode using halide SEs.
During cycling, polycrystalline NCM shows significant particle fractures, which leads to contact loss and cell degradation. However, single-crystalline NCM does not suffer from noticeable cracking. According to a recent paper, polycrystalline NCM demonstrated significantly more oxygen release than single-crystalline NCM. R20 Because the carbonate content was very similar to the NCM employed in this paper and the cells were identical, this result indicates that the different cracking behavior (exposure of fresh and reactive surfaces) is coupled with the release of oxygen.
To probe the gas evolution in situ in ASSB cells employing single-crystalline NCM88 with Li2ZrCl6 and ZrO2-2Li2ZrCl5F, differential electrochemical mass spectrometry (DEMS) measurements were carried out for ASSB cells at 0.1C ( Figure R4). During the operation of ASSB cells (first cycle between 3.0 -4.3 V (vs. Li/Li + ) and second cycle between 3.0-4.8 V (vs. Li/Li + )), no noticeable gas evolution was observed. The outstanding electrochemical oxidative stability of halide SE and the mechanical integrity of the single-crystalline feature of NCM88 suppress the phase transition of CAMs and thus inhibit oxygen release. Response to comment 10: We are thankful to the reviewer for the comment. The oxygen-substituted Li2ZrCl6 (Li12+xZr6Cl36-xOx where x = 1~4) structures were prepared by enumeration technique. We considered all possible Cl and O orderings at Cl sites and selected 50 configurations with the lowest electrostatic energy. All the structures are fully relaxed with DFT calculation. We checked the channel size of the most stable structure at each composition ( Figure R5a), and the optimal point which has the largest channel size and the highest lattice volume is expected to be x = ≈3 in Li12+xZr6Cl36-xOx (Li15Zr6Cl33O3). Those are mainly attributed to the lattice expansion effects by oxygen and the lattice contraction effects by lithium. Moreover, as our EXAFS data show the increased bond length of Zr-Cl in ZrO2-2Li2ZrCl6 HNSE (Figure 1c), we compared the average bond length and radial distribution functions (RDFs) of Zr-Cl in the Li12+xZr6Cl36-xOx structures. The average bond length of Zr-Cl increases significantly at x = 3 or more, which can be also confirmed by the RDF results ( Figure R5b).
Among the generated structures of Li15Zr6Cl33O3 (Li2.5ZrCl5.5O0.5), we selected the structures in which oxygen is placed at one side (since oxygen substitution mainly occurs at the interface of Li2ZrCl6 and ZrO2 and chemical potential difference drives the degree of oxygen substitution gradient). For 20 structures, we conducted short AIMD screening during 20 ps at 700 K with topological analysis and the characteristics of structures that show rapid diffusion were analyzed ( Figure R6). As the repulsion between adjacent ZrCl6-xOx polyhedra increases, the ion conduction channel widens, enabling rapid Li + diffusion. Here, when the oxygen ions of neighboring ZrCl6-xOx polyhedra face each other (agglomerated) in the c-axis direction, such effects are maximized, resulting in fast Li + diffusion. Furthermore, the energy difference between the most stable Li2.5ZrCl5.5O0.5 structure (O-dispersed) and one with such characteristics (O-agglomerated, Figure 3a) is only 9.4 meV/atom. Thus, oxygenagglomerated structures can be sufficiently formed when anion exchange occurs at the interfaces of ZrO2-Li2ZrCl6.
In the revised manuscript, Supplementary Figure 13 and Supplementary texts have been added.  Comment 11. In Fig. 4, the authors claimed that "Li2ZrF6 and Li3Zr4F19 were thermodynamically stable passivating interphases that were decomposed from Li2ZrCl5F." (Line 730-731) (1) I am wondering whether the LiF can be decomposed from Li2ZrCl5F. Why or why not?
Response to comment 11: We are thankful to the reviewer for the insightful comment.
(1) Our calculations show that Li2ZrCl5F (LZCF) can be decomposed into the LiF phase at low voltage (Li insertion) as suggested in Supplementary Table 14. However, since LZCF is used as a catholyte in our system that operates down to 3.0 V (vs. Li/Li + ), it is not expected to form LiF phase.
(2) We summarize their similarities and differences as following; Similarities: By the influence of negative ions (fluorine), those lithium fluoride compounds have higher oxidation stability than other halide compounds such as chlorides and oxides. 21 Our DFT calculations also show that they have very high band gaps (6.31, 5.63, and 8.65 eV for Li2ZrF6, Li3Zr4F19, and LiF, respectively), suggesting that they can be a passivating phase (electronic insulator and Li + conductor) for cathode materials ( Figure R7). Differences: When Li2ZrCl5F is exposed to high-voltage conditions, it will be sequentially decomposed into Li2ZrF6 and Li3Zr4F19 phases, whereas LiF can be formed under very low-voltage conditions. Li-Zr-F ternary compounds (6.526 and 6.613 V for Li2ZrF6 and Li3Zr4F19, respectively) tend to be more stable than LiF (6.332 V) at high voltage ( Figure R7). LiF can also be formed as a reduction-limited decomposition product of Li-Zr-F ternary compounds, and it has a much lower oxidation voltage limit (0 V) than Li2ZrF6 (1.239 V) and Li3Zr4F19 (1.506 V). For the cells tested at an elevated temperature of 60 o C, when the LPSCl monolayer was used so that LPSCl was in direct contact with cathodes, the cathodes using ZrO2-LZCF significantly outperform those using LZC (Figure 5b,c). In contrast, when the LZCF|LPSCl bilayer was used so that LPSCl was not exposed to high voltages, the capacity degradation for using LZC was substantially alleviated but still significant. These results suggest two origins of halide SEs for the degradation of the cathodes using LZC at elevated temperature: i) Incompatibility with LPSCl at high voltages. ii) High-voltage instability.
First, to assess the former, the DFT calculations and control EIS experiments were carried out. The mutual reaction energies of the halide/LPSCl SE mixtures were over 0.4 eV/f.u., strongly indicating their poor compatibility (Supplementary Table 16). For both LZC and ZrO2-LZCF, the EIS experiment using Ti|(halide-LPSCl mixture)|Ti cells stored at 60 ℃ showed marginal changes in Nyquist plots after a week (Supplementary Figure 31), indicating a stable halide-LPSCl interface under no electrochemical driving forces. A control EIS experiment was also performed using (halide-C)|LPSCl|(Li-In) cells charged to 4.3 V (vs. Li/Li + ) at 60 °C so that the halide-LPSCl interfaces are subjected to the high voltage (Supplementary Figure 32). The result exhibited the continuously increased impedance for the cells using LZC. In contrast, the impedance increased for an hour and then saturated for the cell comprised of the ZrO2-LZCF HNSE. This result indicates excellent passivating behaviour for using ZrO2-LZCF. From these results, it can be concluded that the halide/sulfide incompatibility is driven electrochemically at elevated temperatures. Importantly, Fsubstituted chloride SEs are free from this halide/sulfide incompatibility issue.
Second, to assess the high-voltage stability, the CV measurements and DFT calculations were carried out. In the course of preparing the revision, reliable CV results for (halide/C)|halide |LPSCl|(Li-In) cells at 0.1 mV s -1 and 30 o C were obtained and are shown in Figure 4c and Supplementary Table 13. ZrO2-2Li2ZrCl5F exhibited a remarkably smaller integrated anodic current of 1.98 mA V g -1 up to 5.0 V (vs. Li/Li + ), compared to Li2ZrCl6 (2.76 mA V g -1 ). The difference became even larger at the second cycle (0.55 vs. 2.00 mA V g -1 for ZrO2-2Li2ZrCl5F and Li2ZrCl6, respectively). Furthermore, Li2ZrCl6 showed a cathodic peak at ≈3.5 V (vs. Li/Li + ) at the first cycle and they intensified further at the second cycle, which is indicated by an asterisk. It is postulated that the cathodic currents originate from decomposition byproducts formed during the prior positive scan. 53 By contrast, ZrO2-2Li2ZrCl5F exhibited negligible cathodic currents, strongly suggesting its excellent passivating behavior. The DFT results consistently revealed that the oxidative limit of Li2ZrCl5F (4.274 V vs. Li/Li + ) was slightly lower than that of Li2ZrCl6 (4.307 V vs. Li/Li + ), but the formation of desirable F-based passivating interphase materials such as Li2ZrF6 and Li3Zr4F19 can increase the range of the anodic limit (Figure 4d and Supplementary Table 14). 37,38 In conclusion, compared with LZC, the fluorinated HNSE ZrO2-LZCF exhibits remarkably improved sulfide (LPSCl) compatibility and high-voltage stability. These explain the drastic enhancement in cycling performance at the elevated temperature for the LCO electrodes by applying ZrO2-LZCF, compared with those for using LZC.
In the revised manuscript, the relevant discussion has been revised. A reference 53 has also been 21 added.

Comment 4. What is the interacton between F and O in ZrO2-LZCF system?
Response to comment 4: We are thankful to the reviewer for the insightful comment. To investigate the interaction between anions (O/Cl and O/F) in ZrO2-LZC and ZrO2-LZCF systems, we built a slab model of ZrO2(010)-Li2ZrCl5.xF1-x(001) and calculated DFT energies with and without anionic exchange in the interface ( Figure R8). The spontaneous driving force (-0.408 eV) for oxygen-chlorine exchange is confirmed as we have proved experimentally and computationally. In the same way, it is also checked whether the exchange between fluorine from LZCF and oxygen from ZrO2 is energetically favorable or not when fluorine is present at the surface. Since oxygen-fluorine exchange has a stronger driving force (-0.534 eV), it is speculated that oxygen-fluorine exchange can occur in ZrO2-LZCF interfaces.

Response:
We are thankful to the reviewer for a thorough review of our manuscript and the overall positive and constructive comments.

Major Comments:
Comment 1. The vendor and purity of the materials utilized in the synthesis of the halide nanocomposite should be reported.
Response to comment 1: We are thankful to the reviewer for the comment. As shown in the Method section, the information on the materials utilized has been reported.

Comment 2.
In the result section, they wrote "the ZrO2-2Li2ZrCl6 HNSE sample as it is a simple binary system and exhibited a much higher Li+ conductivity of 1.1 mS cm-1 than Li2ZrCl6 (0.40 mS cm-1)" What is the conclusion here? Are they claiming that ZrO2 in between grains modifies the texture of the grain boundary?
Response to comment 2: We are thankful to the reviewer for the comment. The presence of insulating materials in SEs is expected to be disadvantageous. Despite 7.86 vol.% of insulating ZrO2 (based on the theoretical formula), the ZrO2-2Li2ZrCl6 HNSE sample showed a much higher Li + conductivity of 1.1 mS cm -1 than that of Li2ZrCl6 (0.40 mS cm -1 ). We revealed that the anomalous enhancement in Li + conductivity originates from the oxygen-substituted Li2ZrCl6 interphase at the populated interfaces in HNSEs, which was suggested by the DFT calculations results and supported by the synchrotron X-ray and 6 Li MAS-NMR results. Figure 1b should be performed and added to this figure. Lattice constants, coordinates, fractional occupations, R factors, Chi^2, and CIF files should be provided. Presently, the authors just provide the lattice constants extracted from their PDF analysis.

Response to comment 3:
We are thankful to the reviewer for the careful comment. We agree with the reviewer that detailed structure information should be provided. However, Rietveld refinement can only be performed for the XRD data with sufficient peak resolution and intensity. Although we measured XRD in synchrotron (λ = 0.1665 Å), XRD peaks for the ZrO2 phase in ZrO2-2Li2ZrCl6 were not observable. In addition, the Li2ZrCl6 peaks featured a broad shape and weak intensity due to low crystallinity and nanosize (revised Supplementary Figure 3). Therefore, even synchrotron XRD data was insufficient to perform Rietveld refinement, which is why we have adopted pair distribution function (PDF) analysis. PDF is a powerful method for quantitatively analyzing the nanosized and disordered materials since both Bragg and diffuse scattering are measured. The PDF refinements for Li2ZrCl6, ZrO2-2Li2ZrCl6, and nZrO2-2Li2ZrCl6 were performed using a multi-phase model, and the lattice parameters with a reliable factor are tabulated in the revised Supplementary Table 9. Also, we attach CIF files obtained from PDF refinement for providing structural information.
In the revised manuscript, Supplementary Table 9 has been modified as follows. Response to comment 4: We are thankful to the reviewer for the comment. As the reviewer's suggestion, we have added Zr K-edge XANES spectra of ZrCl4 and ZrO2 references in Supplementary Figure 5. The bond strength differs with ligand elements and thus can affect absorption edge energy. As Zr-Cl bonding is weaker than Zr-O bonding, the absorption edge of ZrO2 is positioned at slightly higher energy than the ZrCl4 spectrum. A comparison of the edge position with the reference spectra confirms that the average oxidation state of Zr is tetravalent (4+) in the ZrO2-2Li2ZrCl6 and nZrO2-2Li2ZrCl6 samples.

Comment 5.
High-resolution transmission electron microscopy is known to damage Li2YCl6. I must admit that the micrographs in Figure S3 are not clear, and the diffraction not neat. I wonder whether Li2ZrCl6 derivatives can be equally damaged by the electron beam. The authors should clarify this.

Response to comment 5:
We are thankful to the reviewer for the helpful comment. As pointed out, vulnerable halide SE materials are damaged by electron beams when conventional TEM measurements are performed. In this regard, cryogenic TEM (cryo-TEM), which emerged as a powerful tool for the analysis of electron-beam-sensitive materials, such as biomolecules, Li metal, and solid electrolyte interphases, should be highly promising for halide SE materials. Accordingly, cryo-HRTEM measurements of nZrO2-2Li2ZrCl6 and ZrO2-2Li2ZrCl6 were conducted at cryogenic temperature (−175 °C) without any exposure to ambient air using a double-tilt cryogenic TEM holder. Clear images and corresponding diffraction patterns were acquired by cryo-TEM measurements, from which local nanostructures are visualized (Figures 1e-g). For nZrO2-2Li2ZrCl6, crystalline ZrO2 nanoparticles with diameters ranging from 20 to 30 nm are embedded in glassceramic Li2ZrCl6 matrices (  28 Comment 6. Furthermore, the claim "Interestingly, the HRTEM images of the ZrO2-2Li2ZrCl6 HNSE (Figure 1f, g) showed that ZrO2 formed a percolating network nanostructure with thickness of only a few nanometres." Does not seem supported by the data. If they don't have data backing this statement, this should be removed.
Response to comment 6: We thank the reviewer for the careful comment. As responded in Comment 5, the cyro-HRTEM images with clear diffraction patterns support that nanograin domains of ZrO2 with sizes < 20 nm are dispersed in the glass-ceramic Li2ZrCl6 matrix and ZrO2 formed a local percolating network nanostructure. However, the thickness of the network nanostructure ranges from 10 to 20 nm and it is unclear whether ZrO2 nanodomains form a universal percolating network inside the Li2ZrCl6 matrix.
In the revised manuscript, the relevant description has been revised.

Comment 7.
The statement "To the best of our knowledge, the Li+ HNSE was the first inorganic superionic conductor that exploited the interfacial effect to promote conduction with ionic conductivity reaching 1 mS cm-1." appears a bit over the top. To me this is just the result of ZrO2 modifying the texture of the microstructure, i.e. grain boundary, and hence improving the ion transport. These types of claims should be toned down.
Response to comment 7: We are thankful to the reviewer for the comment. The interfacial conduction effect was first reported by Liang et al. in 1930. 42 Specifically, the Li + conductivity of LiI was improved from 10 -7 to 10 -5 S cm -1 by the addition of metal oxides such as Al2O3. Since then, for various heterostructured materials, anomalous enhancements in ionic conductivity have been reported. 42,43,44,45,46,48 The examples include LiF/silica films (6 × 10 −6 S cm −1 at 50 °C), 45 LiBH4-LiI/Al2O3, (1 × 10 −4 S cm -1 at 20 °C), 48 and PAN with LiClO4/Li0.33La0.557TiO3 (LLTO) nanowires (6.1 × 10 −5 S cm -1 at 30 ∘ C). 46 However, none of these SEs showed ionic conductivities exceeding 1 mS cm -1 at room temperature. To promote the interfacial ionic conductivity effect, a large interfacial area between two different materials is imperative. In most previous literature, the simple mixing of two different materials has been a common practice. This protocol has limitations in terms of downsizing and the high cost of nano-fillers. In this work, we have demonstrated that ZrO2 nanograins with only a few nanometers could be generated in situ by the straightforward mechanochemical method, which is also unprecedented and advantageous in terms of cost. The higher Li + conductivity of the ZrO2-Li2ZrCl6 HNSE (1.3 mS cm -1 ) than that of nZrO2-Li2ZrCl6 prepared by mixing ZrO2 nanoparticles with Li2ZrCl6 (0.60 mS cm -1 ) highlights the novelty of our strategy. Also, considering that conventional compositional tuning, such as ion substitution, has been common to enhance the ionic conductivity of halide SEs, our approach unlocks a new dimension for superionic conduction in materials chemistry. In previous literature regarding interfacial conduction, underlying mechanisms have remained unclear, although various effects, such as the space charge layer and crystallinity effects, were proposed. 42,43,45,47 In our study, we have revealed that the interfacial superionic conduction originates from the oxygen-substituted compounds at the interface. Finally, we demonstrated that our strategy is not restricted to specific compositions. The HNSEs could be formed for Na analogues (0.04 mS cm -1 for ZrO2-Na2ZrCl6, and 0.11 mS cm -1 for 0.13ZrO2-0.61NaCl-0.26Na2ZrCl6) and multi-metal HNSEs such as Al2O3-3Li2ZrCl6 and SnO2-2Li2ZrCl6, which show even higher Li + conductivities (max. 1.6 mS cm -1 for SnO2-2Li2ZrCl6) than ZrO2-2Li2ZrCl6. In summary, our strategy is the first case that achieves the ionic conductivity of 1 mS cm -1 by the interfacial superionic conduction effect. It is also universal and expandable. Moreover, we believe that the elucidation of the mechanism is of 29 significance to the design of superionic conductors. Tables S3 and S4 have a sizeable amount of both ZrO2 and LiCl (or NaCl for the Na analogues) and are far from being phase pure. Now, this study clearly glosses over the importance of these impurities on the texture of grain boundaries, and thus the conductivity properties. A SEM analysis of the most important component is required. There may be a significant statistic of particle size which may also impact these observations.

Comment 8. Furthermore, it seems that many compounds reported in
Response to comment 8: We are thankful to the reviewer for the comment. Based on the DFT calculation results in Supplementary Figure 2, Li2ZrCl6 is the only stable phase in the ZrO2-ZrCl4-LiCl ternary region. Therefore, it is possible to synthesize various HNSEs including LiCl or ZrO2 with possible off-stoichiometry. At the Li2ZrCl6/ZrO2 interface, O-substituted Li2ZrCl6 with high Li + conductivity is formed. The DFT calculation, NMR, and PDF analysis results support the formation and interfacial superionic conduction of the O-substituted Li2ZrCl6 in ZrO2-2Li2ZrCl6. Interestingly, 0.53LiCl-Li2ZrCl6 showed higher ionic conductivities of 0.70 mS cm -1 than that for Li2ZrCl6 (0.40 mS cm -1 ). Even though this work focused on the ZrO2/Li2ZrCl6 interfaces, the enhanced ionic conduction at the LiCl/Li2ZrCl6 interfaces is of course an intriguing issue which could further be investigated. Generally, highly crystalline SE domains can be identified at the grain and grain boundary at the micrometer scale by SEM measurement. However, mechano-chemically prepared HNSEs include poorly crystalline and nanosized grains of ZrO2 and Li2ZrCl6. Thus, the identification of the grain and grain boundary of ZrO2 and Li2ZrCl6 by SEM is not possible due to the limited resolution of SEM. Therefore, we have carried out the cryo-HRTEM analysis for HNSEs. The cryo-HRTEM analysis results of ZrO2-2Li2ZrCl6 demonstrate that poor-crystalline nanograins of ZrO2 are spread randomly in the glass-ceramic Li2ZrCl6 matrix (Figures 1e-g and Supplementary Figure 6). The significantly larger-area percolating interfaces for ZrO2-2Li2ZrCl6 compared with nZrO2-2Li2ZrCl6 are confirmed by this observation. Please read our response to the Comment 5 for the detail. Nevertheless, according to the reviewer's comment, we obtained SEM images for Li2ZrCl6, ZrO2-2Li2ZrCl6, and ZrO2-2Li2ZrCl5F ( Figure R9). The samples showed irregular particles with sizes at the micrometer scale without distinct differences from each other. Comment 9. Figure 2d, it's unclear why there is just a datapoint for nZrO2-2Li2ZrCl6. Why couldn't the author measure the ion conductivity at lower or higher temperatures?
Response to comment 9: We are thankful to the reviewer for the comment. According to the comment, ionic conductivities of nZrO2-2Li2ZrCl6 were obtained at different temperatures ( Figure  2c). In the whole temperature range, Li + conductivities of nZrO2-2Li2ZrCl6 were lower than those of HNSE (ZrO2-2Li2ZrCl6) but higher than those of LZC.
In the revised manuscript, the corresponding data have been added in Figure 2c.  Response to comment 10: We are thankful to the reviewer for the insightful comment. All the EIS data for the conductivity measurements were fitted using an equivalent circuit model shown in Figure  R10a. The equivalent circuit model and the fitted results are also provided in Supplementary Figure  10 and Supplementary Table 4. As pointed out, in Nyquist plots of SE materials, signals for grain (or bulk) and grain-boundary resistances evolve at higher and lower frequencies, respectively, and could often be deconvoluted especially at low temperatures. R21-23 In SEs, the grain and grain boundary capacitance values usually lie in ≈10 −12 F and 10 −11 -10 −8 F, respectively. R22,23 However, the EIS data for ionic conductivity in the original manuscript do not show two distinct semicircles or asymmetric features. Therefore, an EIS measurement for ZrO2-2Li2ZrCl6 was conducted at a lower temperature of −40 °C. The corresponding Nyquist plot is presented in Figure R11. Unfortunately, it also showed one semicircle with decent symmetry. Nevertheless, the data for ZrO2-2Li2ZrCl6 were fitted with an equivalent circuit model representing the grain and grain boundary contributions as shown in Figure  R10b. However, it fails to deconvolute the contributions. Instead, the data were fitted well with the equivalent circuit model with the one parallel-connected RC component ( Figure R10a). And the capacitance value for the single semicircle is 27.71 pF, which corresponds to the bulk (grain) transport. Therefore, the contribution for total resistances for HNSEs by grain boundary resistances is likely marginal.  Comment 11. The authors claimed that "Li+ diffusion was ~11 times faster for LZCO than LZC at 300 K". This is a very big claim considering that AIMD at a temperature as low as 300 cannot be considered converged. Perhaps is more valuable to make this comparison at higher temperatures.
Response to comment 11: As the reviewer knows, an extremely long computational time for AIMD simulations is required to predict statistically reliable Li + diffusivity at room temperature (300 K) due to slow Li + hopping at low temperatures. Thus, it is technically impossible to directly predict the diffusivity of material with a conductivity of about 1 mS cm -1 at 300 K using AIMD. This is why Li + diffusivity at 300 K has been predicted by extrapolating a few values of Li + diffusivity calculated above 500 K, and we also followed the general method. R24-26 However, we agree with the reviewer that this is a very big claim as we assume that the diffusivity follows the Arrhenius relationship without the slope change from high to low temperatures. As the reviewer suggested, we will present diffusivity data at high temperatures as an additional supplementary table to prove not only that it is faster for LZCO than LZC at high temperatures but also that the slope is less steep for LZCO, which reflects higher diffusion at 300 K. In addition, the statements have been revised to avoid misunderstanding of the results.
In the revised manuscript, Supplementary Table 12 has been added, and the relevant discussion has been revised as follows.